Î²-FeOOH: An Earth-Abundant High-Capacity Negative Electrode

Page 1 of 11

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Chemistry of Materials

β-‐FeOOH: An Earth-‐Abundant High-‐Capacity Negative Electrode Material for Sodium-‐Ion Batteries Linghui Yu,† Luyuan Paul Wang,†, ‡ Shibo Xi,§ Ping Yang,⊥ Yonghua Du,§,* Madhavi Sriniva-‐ san,†,§ and Zhichuan J. Xu†,§,* †

School of Materials Science and Engineering, Nanyang Technological University, 50 Nanyang Avenue, 639798, Singapore §

Energy Research Institute@NTU, Nanyang Technological University, 50 Nanyang Drive, 639798, Singapore

‡

Institute of Chemical and Engineering Sciences, A*STAR, 1 Pesek Road, 627833, Singapore

⊥ Singapore Synchrotron Light Source (SSLS), National University of Singapore, 5 Research Link, 117603, Singapore.

ABSTRACT: Thanks to the great earth abundance of sodium and decent energy densities, sodium-‐ion batteries are prom-‐ ising alternative energy storage devices for large-‐scale applications. Developing cheap, safe and high-‐capacity sodium-‐ion battery anode materials is one of critical challenges in this field. Here, we show that β-‐FeOOH is a very promising low-‐ cost anode material, with a high reversible capacity (>500 mAh g-‐1 during initial cycles). The fundamental characteristics associated with the discharge/charge processes, in terms of the redox reactions, formation/deformation of the solid elec-‐ trolyte interface (SEI) layers, and structural and morphological changes, are comprehensively investigated. In addition, a comparison study shows that the smaller-‐sized FeOOH has more serious kinetic restrictions, and thus lower capacities, while it shows better cyclability than the bigger one. Origins of the large overpotential are discussed and it is suggested that the overpotential should be mainly due to the features of the surface concentration dependent potential and the slow diffusion of Na+; and the presence of the SEI layers may also contribute to the overpotential.

INTRODUCTION Sodium (Na)-‐ion batteries (NIBs) have been receiving increasing research attention in recent years, due to the potential use for large-‐scale applications, such as station-‐ ary energy storage.1-‐5 Compared with lithium-‐ion batteries (LIBs), NIBs may have lower but still considerably high energy densities. However, the great abundance of sodi-‐ um makes NIBs potentially much cheaper, hence, very attractive for large-‐scale applications where both cost and energy density are key factors. The development of NIBs has been snagged perhaps mainly by the lack of suitable electrode materials. Alt-‐ hough NIBs are conceptually identical to LIBs, the large radius of Na+ (0.3 Å larger than that of Li+) is an intrinsic drawback that results in sluggish diffusion kinetics of Na+ in electrode materials, leading to lower capacities or even electrochemical inactivity.1, 6-‐9 Intensive efforts have been devoted to improving and exploring both cathodes and anodes. For anodes, hard carbons are early reported ma-‐ terials with significant sodium storage capacities at low potentials vs. Na+/Na.10-‐14 So far, much progress has been made in terms of capacity and cycalibility. Even so, they seems to approach a limiting value around 300 mAh g-‐1. Titanium-‐based materials, such as, Na2Ti3O7,15-‐18 Li4Ti5O12,19-‐23 Na0.66[Li0.22Ti0.78]O2,24 TiO2,25-‐28 and Na-‐ Ti2(PO4)3,29-‐31 have emerged as a promising class of anodes for NIBs. Although they have a relatively low capacity

between 100 – 300 mAh g-‐1, several of them have shown excellent cylability and have been proposed for long-‐term usage. Research work on high-‐capacity anode materials focus-‐ es mainly on tin-‐ and antimony-‐based materials, such as, Sn,32-‐35 Sb,36-‐40, SnSb,41 SnO2,42, 43 SnS,7, 44 SnS2,7, 45, 46 Sb2O4,47 and Sb2S3.48 With the electrochemical formation of Na-‐Sn or Na-‐Sb alloys, they own high theoretical ca-‐ pacities, for instance, 847 and 660 mAh g-‐1 for pure tin and antimony, respectively. Many reports have shown that a large proportion of the theoretical capacities could be realized, from 500 mAh g-‐1 up to close to the theoreti-‐ cal capacities. Despite the high capacity, these materials usually exhibit poor cyclability and there are concerns on the costs of NIBs raised by using relatively expensive tin or antimony. Their earth abundance may also fail to meet large-‐scale grid applications. Phosphorus-‐based materials, such as P,49-‐51 Sn4P3,52-‐54 and Ni3P,55 are another class of high-‐capacity anodes reported for NIBs. For pure phos-‐ phorus, it has the highest theoretical capacity (2596 mAh g-‐1) of all the anode materials reported. And a capacity up to 1800 – 1900 mAh g-‐1 has been achieved by several groups, though fades fast upon cycling. However, the use of phosphorus-‐base materials involves formation of highly toxic Na3P,56-‐58 which additionally increases safety haz-‐ ards.

ACS Paragon Plus Environment


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

While the development of high-‐capacity alloying and phosphorus-‐based materials is accompanied by cost and/or safety considerations, the use of iron-‐based mate-‐ rials becomes highly attractive, because of their low costs, environmental friendliness, and high theoretical capaci-‐ ties. Iron-‐based materials, such as Fe2O3 and Fe3O4, have been widely investigated for LIBs.59, 60 They can electro-‐ chemically react with lithium to form metallic iron and deliver high theoretical capacities, e.g., 1005 and 926 mAh g-‐1 for Fe2O3 and Fe3O4, respectively. Similar mechanisms of formation of metallic iron have been also proposed for their sodium storage.61-‐63 But the sodium storage capaci-‐ ties are usually limited by the slow kinetics associated with the large radius of Na+. Among the reported materi-‐ als, several relatively high capacities were obtained based on nano-‐sized particles in composites with carbons. For example, a Fe3O4 with particle size less than 10 nm in a matrix of carbon showed a capacity of ~350 mAh g-‐1 at a low current density.64 Fe2O3 nanocrystals anchored on graphene could have a stable capacity around 400 mAh g-‐1 at 100 mA g-‐1.65 A 5 nm γ-‐Fe2O3 embedded in a matrix of carbon exhibited a high capacity that could be stabilized at ~700 mAh g-‐1 at 100 mA g-‐1.66 So far, in addition to ki-‐ netic problems, iron-‐based materials always show very low initial reversibility. Most of them have an initial Cou-‐ lombic efficiencies (CEs) of lower than 60%,62, 64, 65, 67, 68 while some of them have a value close to 70%.63 Such low CEs are unacceptable for practical applications. Herein, we present a promising sodium storage anode β-‐FeOOH with a high reversible capacity (627 and 523 mAh g-‐1 for the first and second cycle, respectively, at a current of 80 mAh g-‐1). The sodium storage processes are investigated by a combination of electrochemistry, X-‐ray absorption near edge structure (XANES) spectroscopy, X-‐ ray diffraction (XRD) and electron microscopy to under-‐ stand the redox reactions, formation/deformation behav-‐ iors of the solid electrolyte interface (SEI) layers, and structural and morphological changes. A comparison study on two different sized materials (rod-‐shape, diame-‐ ter 5 vs. 53 nm) is also carried out and preliminary results show interesting size-‐dependent properties, that the 5 nm β-‐FeOOH has more serious kinetic restrictions, and thus lower capacities, while it shows better cycling stability. Moreover, while the origins of the large overpotential of electrode materials have been usually considered to be complicated and difficult to determine,60, 69 we suggest, based on our study, that the overpotential should be mainly due to the features of the surface concentration dependent potential and the slow diffusion of Na+; the presence of the SEI layers may also contribute to the overpotential. EXPERIMENTAL SECTION Synthesis. A simple hydrolysis method was used to con-‐ trol the size of β-‐FeOOH nanorods.59, 70 For the synthesis of β-‐FeOOH nanorods with a mean diameter of 53 nm (see Supporting Information Figure S1 for size distribu-‐ tion), 10 mL of 0.5 M aqueous FeCl3 (Sigma-‐Aldrich) was

Page 2 of 11

added into 70 mL of H2O in a 100 mL plastic bottle with a 0.8 mm needle hole. After shaking for 1–2 min, the bottle was kept in an oven at 100 °C for 24 h. The precipitates were collected after cooling down and washed with de-‐ ionized water and ethanol several times and finally dried at 80 °C overnight. For the synthesis of β-‐FeOOH nano-‐ rods with a mean diameter of 5 nm, 5 mmol of LiOH (Alfa Aesar) was added in the beginning while all other condi-‐ tions were kept the same. Characterization. XRD analysis was carried out with a Shimadzu thin-‐film diffractometer using Cu Kα radiation for both the prepared materials and the electrodes after discharge/charge. Field emission scanning electron mi-‐ croscopy (FESEM) images were collected on a JEOL JSM-‐ 6340F microscope. High-‐resolution transmission electron microscopy (HRTEM) images were observed using a JEOL JEM 2010 microscope. XANES was carried out at Singa-‐ pore Synchrotron Light Source (SSLS), XAFCA beamline. For the XRD and XANES analysis on the material after discharge/charge, all the electrodes were sealed with Kap-‐ ton tape to prevent oxidation. Electrochemical Measurements. The active material β-‐ FeOOH nanorods were mixed with graphite (Sigma-‐ Aldrich), conductive carbon (Super P) and binder (sodi-‐ um carboxymethyl cellulose, CMC, Sigma-‐Aldrich) in de-‐ ionized water in a weight ratio of 50 : 30 : 5 : 15. The re-‐ sultant slurry was coated onto a copper current collector and dried at 80 °C overnight under vacuum. The graphite was used to improve stability of the electrode.59 The elec-‐ trodes were assembled with Li metal as counter elec-‐ trodes in a 2032 coin cell in an Ar-‐filled glove box. Glass fiber separators (EL-‐CELL, ECC1-‐01-‐0012-‐C) were used to keep the electrodes apart. The electrolyte was 1 M NaClO4 (Sigma-‐Aldrich) in propylene carbonate (Sigma-‐Aldrich) with 5 wt% fluoroethylene carbonate (FEC, Sigma-‐ Aldrich). Galvanostatic discharge and charge (0.005 – 2.8 V) tests were performed on Battery Testing Equipment (Neware Electronic Co., China) at different current densi-‐ ties under ambient temperature. Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) meas-‐ urements were performed on a Bio-‐Logic SP-‐150 potenti-‐ ostat. The frequency range and AC amplitude used for the EIS measurements were 1 M Hz – 0.01 Hz and 10 mV, re-‐ spectively. RESULTS AND DISCUSSION The β-‐FeOOH studied here was synthesized by a simple hydrolysis process. They own rod-‐like shapes with con-‐ trolled sizes (mean diameters: 53 and 5 nm, Figure S1). The phase structure was confirmed by XRD analysis (Fig-‐ ure S2). The specific surface areas are 16 and 105 m2 g-‐1 for the 53 and 5 nm materials, respectively.70 We first investi-‐ gate the sodium storage behaviors of the material based on β-‐FeOOH rods with mean diameter of 53 nm. And then conduct a comparison study with a 5 nm β-‐FeOOH. Figure 1a shows the first five discharge/charge curves of the 53 nm β-‐FeOOH at a current density of 80 mA g-‐1. The first discharge (sodium insertion) curve has one short


plateau between 1.1 and 1.2 V vs. Na+/Na at the beginning of the sodium insertion process, followed by a falling po-‐ tential till the end of discharge at a potential of 0.005 V vs. Na+/Na. There is a phase transformation process as the structure loses long-‐range order without any significant XRD peaks after insertion of sodium (Figure S3). Upon charge (sodium extraction), the voltage profile is lack of a well-‐defined plateau, composed of only sloping regions, indicating the poor reversibility of the reaction occurred at the plateau during the initial discharge. The phase transformation behavior is found to be irreversible as the material remains amorphous after charged back to 2.8 V. The discharge and charge processes deliver a capacity of

2.5 2.0 1.5 1.0 0.5 0.0 0

200 400 600 800 1000 Capacity (mAh g-1)

792 and 627 mAh g-‐1, respectively, corresponding to an irreversible loss of 165 mAh g-‐1 and a CE of 79%. Such a high reversible capacity is close to the theoretical capacity of alloying Sb anode. The CE is at least 10% higher than those reported previously for iron-‐based materials. 62-‐65, 67, 68 To be able to exclude the capacity contribution of the carbons (Super P and graphite) which were used to im-‐ prove the conductivity and stability of the electrode,59 we also tested their sodium storage capacity, and found a very low reversible capacity of ~20 mAh g-‐1 (Figure S4), which is negligible compared to the large capacity deliv-‐ ered from β-‐FeOOH.

b

100 900

Discharge Charge

600

85

75 2

4 6 8 Cycle number

10

0.4 0.2

90

80

300 0

c

95

CE

CE (%)

1st 2nd 3rd 4th 5th

3.0

Capacity (mAh g-1)

a Potential (V, vs. Na+/Na)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Current (A g-1

Page 3 of 11

0.0 -0.2 -0.4 -0.6

sh Fre 1st 2nd 3rd 4th

0.0 0.5 1.0 1.5 2.0 2.5 3.0 + Potential (V vs. Na /Na)

Figure 1. Electrochemical properties of β-‐FeOOH (Dave = 53 nm). a) Galvanostatic discharge/charge curves at a current density of 80 mA g-‐1; b) first 10 cycle capacity and CE; c) CV curves at a scanning rate of 0.2 mV s-‐1.

In the second cycle, both of the discharge and charge profiles are quite unlike the first. There is also an absence of a well-‐defined plateau during the second discharge and the sloping profile features a significant voltage step from 0.4 to 0.3 V. During charge, the profile becomes smoother compared with the first cycle. These suggest that the ma-‐ terial is instable after one cycle. During the third cycle, the discharge profile is still distinctively different from the second, while the charge profile remains nearly the same. In the subsequent cycles, the discharge and charge profiles of each cycle are almost invariant, and only shows a small but continuous reduction in capacity. Because the shape of the curves start to be almost the same from the second charge (i.e. after the second discharge), the mate-‐ rial become relatively stable since. The discharge and charge capacities of the second cycle are 580 and 523 mAh g-‐1, respectively. The capacity decreases gradually after the second cycle. A charge capacity of 351 mAh g-‐1 is obtained after 10 cycles (Figure 1b). The CE reaches close to 100% after a few cycles. The material also owns a good rate ca-‐ pability, at a current of 500 mA g-‐1, it has a reversible ca-‐ pacity of 509, 456 and 187 mAh g-‐1 for the 1st, 2nd and 100th cycle, respectively (Figure S5). CV was also employed to investigate the sodium storage behavior of β-‐FeOOH. Figure 1c shows CV profiles of the 53 nm β-‐FeOOH. Each cycle presents several peaks during both reduction and oxidation, characteristic of multiple-‐ intermediate processes of insertion and extraction. Dur-‐ ing the first cycle, the material show reduction peaks at

0.90, 0.59 and near 0 V; and one pronounced oxidation peak at 1.39 V associated with several shoulder peaks on both sides. During the second cycle, there is a significant reduction in intensity of the main peaks because the reac-‐ tions are partially reversible. The reduction peaks are al-‐ most at the same positions with the first cycle, while the main oxidation peaks appear at different positions due to an activation process. From the second oxidation cycle, the characteristics of the CV profile remain practically identical, with only small reduction in peak intensity and small shifts of the peak positions, indicating the material becomes more stable, consistent with the discharge and charge processes. The peak at 0.90 V can be assigned to the reaction at the short plateau during the first discharge corresponding to the first iron reduction step. The most pronounced reduction peak at 0.59 V during the first cy-‐ cle corresponding to the second iron reduction step. It reduces fast in the following cycles, and nearly vanishes in the fourth cycle. The reduction current near 0 V is from a combination of reduction of iron and formation of SEI layers as confirmed by XANES and microscopy studies below. The oxidation peaks during anodic cycles originate from deformation of the SEI layers and corresponding multiple-‐intermediate sodium extraction reactions. To further understand the redox processes associated with sodium insertion and extraction, we carried out XANES studies on the electrodes at different sodiation (discharge/charge) states during the first two cycles. Comparison of the Fe K-‐edge XANES spectra allows us to


Chemistry of Materials acquire valence information of iron in the electrode as the overall edge position shifts to lower energies upon reduc-‐ tion, due to the decrease in binding energy of the Fe 1s2 electrons. The spectra are shown in Figure 2a – f, the points chosen for analysis shown in Figure 2g, and the corresponding 1st derivative plots shown in Figure S6. During the first discharge (Figure 2a), we can see the ab-‐ sorption edge position shifts continuously till the cut-‐off potential of 0.005 V vs. Na+/Na. This means the reduction of iron occurs over the whole potential region. It is ob-‐ served that the spectra for the low potential electrodes (0.3 and 0.005 V) exhibit a significantly higher pre-‐edge jump compared with those at high potentials, indicating the presence of metallic iron in the electrodes at low po-‐

e

Fe 1D, 0.005 V 1C, 0.55 V 1C, 2.0 V 1C, 2.8V β-FeOOH

7110 7125 7140 7155 Energy (eV)

f

1D, 0.005 V 2D, 0.005 V

7140

7200 7260 Energy (eV)

d

β-FeOOH 1C, 2.8 V 2D, 0.8 V 2D, 0.3 V 2D, 0.005 V Fe

7110 7125 7140 7155 Energy (eV)

g

1D

4.0

β-FeOOH 1D, 1.1 V 1C, 2.8 V 2C, 2.8 V

7110 7125 7140 7155 Energy (eV)

Normalized absorption (a.u.)

c Normalized absorption (a.u.)

7110 7125 7140 7155 Energy (eV)

tentials, which is further confirmed on the extended X-‐ray absorption fine structure (EXAFS) spectrum with the presence of the iron peak (Figure S7). However, even at 0.005 V, the absorption edge position is clearly higher than that of the iron reference, indicating the average valence of iron at 0.005 V is above 0. These results demonstrate the formation of metallic iron at low poten-‐ tials, but only partial of the iron in the material is reduced to the metallic state, which is consistent with that the 1st discharge capacity of 792 mAh g-‐1 is lower than the theo-‐ retical capacity of 905 mAh g-‐1 from Fe(III) to Fe(0). This partial reduction of iron is different from its Li storage behavior where the iron could be completely reduced to its lowest valence.59

Potential (V, vs. Na+/Na)

β-FeOOH 1D, 1.1 V 1D, 0.8 V 1D, 0.3 V 1D, 0.005V Fe


b



a


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 4 of 11

Fe 2D, 0.005 V 2C, 0.55 V 2C, 2.0 V 2C, 2.8 V β-FeOOH

7110 7125 7140 7155 Energy (eV) 2D

1C

3.2

2.8

2.4

2.0

1.6

1.1

0.8

0.8 0.3

0.8

0.55 0.005

0.0 0

10

20 Time (h)

2C 2.8 2.0 0.55 0.3 0.005 30

Figure 2. Fe K-‐edge XANES spectra for β-‐FeOOH (Dave = 53 nm) at different sodiation states. a) Electrodes after the first discharge to 1.1, 0.8, 0.3 and 0.005 V, comparing the β-‐FeOOH and metallic Fe as references; b) electrodes after the first discharge to 0.005 V and after the first charge to 0.55, 2.0 and 2.8 V, and references; c) electrodes after the first charge to 2.8 V and after the second discharge to 0.8, 0.3, and 0.005 V, and references; d) electrodes after the second discharge to 0.005 V and after the second charge to 0.55, 2.0 and 2.8 V, and references; e) electrodes after the first and second charge to 0.005 V; f) electrodes after the first dis-‐ charge to 1.1 and after the first and second charge to 2.8 V, and β-‐FeOOH reference; g) points chosen for XANES analysis. 1D, 1C, 2D, and 2C in the legends stand for first discharge, first charge, second discharge, and second charge, respectively.

Upon subsequent charge (Figure 2b), oxidation of iron starts above 0.55 V as the plots almost overlaps at 0.005 and 0.55 V. At the final charge potential of 2.8 V, iron is still not fully oxidized to Fe(III) as the absorption edge position is located slightly lower than the one for the pris-‐ tine β-‐FeOOH. This results in an irreversible loss of the capacity consistent with the galvanostatic cycling meas-‐ urement. To determine the oxidation state at 2.8 V after one cycle, we compare the spectrum at this potential with the one at 1.1 V during the first charge, and find the ab-‐ sorption edge position at 2.8 V is higher than the one at 1.1 V. The valence of iron at 1.1 V is +2.6 calculated from the capacity of 128 mAh g-‐1 at this potential, assuming

there is no formation of SEI layers at this relatively high potential and all the capacity is delivered from the reduc-‐ tion of iron. Thus, the valence of iron at 2.8 V is between +2.6 and +3. The irreversible loss of the capacity caused by the incomplete oxidation of iron should be less than 128 mAh g-‐1, lower than the total loss of 165 mAh g-‐1 in the first cycle, indicating there are other irreversible (or par-‐ tially irreversible) reactions, which can be attributed to the formation/deformation of SEI layers as confirmed below by field emission scanning electron microscopy (FESEM) and high-‐resolution transmission electron mi-‐ croscopy (HRTEM).


Page 5 of 11

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


The spectra for the second cycle show the same trend (Figure 2c and d). Upon discharge, iron is gradually re-‐ duced until 0.005 V; and its oxidation starts above 0.55 V during charge. The oxidation state at 0.005 V is very simi-‐ lar to the one after the first discharge with their edge po-‐ sition located at quite the same position (Figure 2e and Figure S6e). However, we can find the two spectra have different shapes, indicating different local chemical envi-‐ ronments of iron. The oxidation state of iron at 2.8 V after two cycles is close to the one at 1.1 V during the first dis-‐ charge (Figure 2f), lower than that after the first charge at 2.8 V, in agreement with the capacity reduction observed. The morphological changes upon sodium insertion and extraction as well as the formation/deformation behaviors of SEI layers were studied by both FESEM and HRTEM. Figure 3a and Figure S8a show the FESEM images of fresh electrode. The original rod-‐like shape of the materi-‐ al remains in the electrode. We can also find small Super P carbon and micro-‐sized graphite used in the electrode in the FESEM images. After insertion of a small amount of sodium at 1.1 V, the material starts to crack into small na-‐ noparticles (Figure 3b), indicating large volume strains upon sodium insertion. After discharged to 0.8 V, there is further cracking of the particles which become signifi-‐ cantly smaller than at 1.1 V (Figure 3c). Discharge to 0.3 V doesn’t lead to an obvious size change of small nanoparti-‐ cles (Figure 3d). However, at this potential, we can easily find large microsized particles (indicated by the ellipse)

with morphology similar to the final discharge state at 0.005 V, which seem to be agglomerates from small parti-‐ cles. At 0.005 V, microsized particles are the domain characteristic of the electrode. There are also a lot of small nanoparticles on the surface (see also Figure S8b). We attribute the appearance of the microsized particles observed at low potentials to the formation of SEI layers. The SEI layers form on the surface of the small particles and make them interconnect with each other. If the layers are thick enough, the boundaries between nanoparticles disappear and they grow into large particles. The SEI lay-‐ ers are formed inhomogeneously, evidenced by the coex-‐ istence of the nano-‐ and micro-‐particles at the low poten-‐ tials. After charged back to 2.8 V, the majority of the SEI layers deforms, and the material remains to be small nano-‐spheres instead of the original rod shape with larger sizes. The cracking of the original particle and formation of the SEI layers are confirmed by HRTEM analysis. At 0.005 V (Figure 3g), particles with size of ~100 nm are observed in SEI matrix. Interestingly, these particles are composed of very small nanoparticles. This is similar to conversion anodes after lithiation.60 At 2.8 V (Figure 3h), we can still find a thin SEI layer on the particles, indicat-‐ ing the deformation of SEI layer is partially reversible. The extreme volume strains observed here has the serious drawback that causes instability of the electrode leading to poor cyclability (Figure 1b and Figure S5).

Figure 3. SEM and TEM images of β-‐FeOOH (Dave = 53 nm) at different sodiation states. SEM images of a) fresh electrode; b, c, d and e) electrodes after the first discharge to 1.1, 0.8, 0.3 and 0.005, respectively; f) electrode after the first charge to 2.8 V. TEM images of the electrodes after the first discharge to 0.005 V (g); and after the first charge to 2.8 V (h). The ellipse in (d) indicates similar morphology observed in (e). The arrows in (g) and (h) indicate the SEI layers.

Finally, we compare the sodium storage performance of the 53 nm material with the 5 nm one. Figure 4 shows the electrochemical properties of the 5 nm β-‐FeOOH. The materials has a relatively low capacity and rate capability

(495 and 266 mAh g-‐1 at 80 and 500 mA g-‐1), while exhibits rather stable cycling behavior. After two cycles, the dis-‐ charge/charge curves overlap at a current of 80 mA g-‐1 and the reversible capacity remains almost the same at each



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

termine the origins of overpotential,60, 69 it seems clear here that the overpotential of the β-‐FeOOH electrode is at least mainly derived from the features of the surface concentration dependent potential and the slow diffusion of Na+.

rate. The improved cycling performance is probably be-‐ cause that smaller particles provide possibility of better contacts between electrode material, conductive carbon and binder which is beneficial for the maintenance of the stability of the electrode, provided that there is sufficient conductive carbon and binder used.

Comparing the two materials in the relaxation meas-‐ urements, we can find that Na+ diffuses slower in the 5 nm material as it takes longer to achieve the quasi-‐ equilibrium state and the 5 nm material also has a large overpotential, for both the fresh electrodes and the ones after three cycles. It’s interesting that the smaller-‐sized material needs longer time to achieve the quasi-‐ equilibrium state because it has a shorter distance for Na+ to diffuse to the very inner part. The slower diffusion of the 5 nm material should lead to the larger overpotential observed. Understanding this difference between the two materials should be helpful in designing materials that can avoid the large overpotential. This would be worth further research efforts. When comparing the measure-‐ ments on the fresh electrodes and on the ones after three cycles, we can find the electrodes after three cycles need longer time to achieve quasi-‐equilibrium states for both two materials, indicating slower diffusion kinetics after cycling. This might be attributed to the slow diffusion of Na+ in the SEI matrix that forms at low potentials and remains still partially at high potentials as discussed above. In this regard, building advanced SEI layers where Na+ can transfer fast will be favorable for reducing the overpotential. This might be achieved by changing elec-‐ trolyte and/or adding additives. It is also found that the overpotentials obtained using the two electrodes after three cycles are larger than the ones before cycling, indi-‐ cating that the presence of SEI contributes to the overpo-‐ tential.

Potential relaxation measurements (Figure 5) per-‐ formed on the two β-‐FeOOH provide further insights into their sodium storage behaviors. Two electrodes at differ-‐ ent sodiation states were used for the measurements for both materials, fresh electrodes (Figure 5a) and electrodes after three cycles (Figure 5b). Sodium was inserted step by step using a very low current of 20 mA g-‐1 for 1 h and wait-‐ ing after each step for 10 h until a capacity of 100 mAh g-‐1 was achieved. Afterwards, the batteries were left until stable, i.e., achieving their quasi-‐equilibrium state. The quasi-‐equilibrium potentials after insertion of a capacity of 100 mAh g-‐1 are 2.64 and 2.61 V for the fresh electrodes with the 53 and 5 nm materials, respectively. Materials with different sizes have different specific surface areas and they are in some cases considered to affect the equi-‐ librium potential of electrode materials due to the surface energy contribution.60, 71 Here, the slight difference, 0.03 V, of the potentials indicates that the surface energy con-‐ tribution has limited effects on the equilibrium potential for β-‐FeOOH for sodium storage. When using the elec-‐ trodes after three cycles, the quasi-‐equilibrium potentials are 2.68 and 2.60 V for the 53 and 5 nm materials, respec-‐ tively. Apart from the possible size difference of these two materials after cracking, the enlarged potential difference, 0.08 V, may be also due to the different sodiation states of the two materials, because they may contain different amount of Na+ after three cycles. In all the cases, it takes very long for the material to achieve the quasi-‐equilibrium state, more than 300 h, being the fastest, for the 53 nm material initially in the fresh electrode, and more than 1000 h or even more than 3000 h for the others. This suggests a very slow Na+ diffu-‐ sion rate in the material. The overpotentials, as indicated in Figure 5, are also found to be very large, >1 V, for all the cases. The slow diffusion and the large overpotentials suggest that the potential of the material highly depend on the concentration of Na+ in the surface part of the ma-‐ terial (similar to a pseudo-‐capacitive behavior72). With the decrease of the surface concentration via diffusion of Na+ toward inner where its concentration is lower, the open circuit potential increases until the equilibrium state is reached. While it is usually thought to be difficult to de-‐

Page 6 of 11

As mentioned, the 5 nm material has slower diffusion of Na+. It suffers from more serious kinetic restrictions dur-‐ ing sodium insertion and extraction, and thus leads to a lower capacity and rate capability. One may also consider other kinetic factors for the relatively low capacity and rate capability of the 5 nm material, such as the conduc-‐ tivity and the interfacial resistance. However, electro-‐ chemical impedance spectroscopy analysis (Figure S9) has shown the two batteries have the same ohmic re-‐ sistance and the 5 nm material has smaller interfacial re-‐ sistances than the 53 nm one.


0.5

0m A g -1 20

00

300

mA g -1

0m A g -1

1.0

10

1.5

600

50

2.0

Discharge Charge

mA g -1 0m A g -1

2.5

b

80

1st 2nd 3rd 4th 5th

3.0

20


a

0.0 0

200 400 600 Capacity (mAh g-1)

800

0

10

20 30 40 Cycle number

50

Figure 4. Electrochemical properties of β-‐FeOOH (Dave = 5 nm). a) Galvanostatic discharge/charge curves at a current density of 80 mA g-‐1; b) cycling performance and rate capability.


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Capacity (mAh g-1)

Page 7 of 11

3.0 a) 2.64 V 2.61 V 2.5 2.0 1.32 V 1.44 V 1.5 1.0 0 500 1000 1500 Time (h) 3.0 b) 2.5 2.0 1.5 1.0 0

-1

100 mAh g 53 nm 5 nm 2000

0

2.68 V 2.60 V 1.66 V 1500

1.67 V

3000 Time (h)

Rest

20 40 Time (h)

60

Rest 100 mAh g-1 Rest 53 nm 5 nm

4500

6000 0

40 80 Time (h)

Figure 5. Time for achieving quasi-‐equilibrium states after insertion of sodium capacity of 100 mAh g-‐1. Measurements performed on a) fresh electrode and b) electrodes after three cycles and when the OCVs became relatively stable. Sodium was inserted step by step using a very low current of 20 mA g-‐1 for 1 h and waiting after each step for 10 h until a capacity of 100 mAh g-‐1 was achieved. Afterwards, the batteries were left until quasi-‐equilibrium. The quasi-‐equilibrium potentials and the overpotentials after insertion of a capacity of 100 mAh g-‐1 of Na+ have been indicated.

CONCLUSIONS Earth-‐abundant β-‐FeOOH has been investigated as a novel anode for NIBs. The 53 nm sample exhibits a high reversible capacity (627 and 523 mAh g-‐1 for the first and second cycle, respectively, at a current of 80 mAh g-‐1. Alt-‐ hough it also shares the same drawback, fast capacity fad-‐ ing, of other high-‐capacity anodes, the material holds great promise for sustainable NIB applications. A range of techniques including CV, XANES, XRD and microscopy were used to understand the dis-‐ charge/charge processes. The material has a phase trans-‐ formation from crystalline to amorphous after insertion of sodium. The phase transformation behavior is irreversi-‐ ble. The sodium insertion process is accompanied by ob-‐ vious particle cracking. The conversion of the material to metallic iron at low potentials is demonstrated. However, the conversion process is incomplete even at a cut-‐off potential of 0.005 V vs. Na+/Na, only partial of the materi-‐

al is converted to metallic iron. At a charge cut-‐off poten-‐ tial of 2.8 V vs. Na+/Na, iron can be oxidized up to a va-‐ lence between +2.6 and +3, instead of back to +3. This contributes to the irreversible loss of the capacity. The formation of the SEI layers at low potentials was also demonstrated. Its deformation occurs during charge pro-‐ cess, but still incomplete at 2.8 V, which also contributes to the irreversible capacity. A comparison of sodium storage properties on two dif-‐ ferent sized materials (5 vs. 53 nm) shows interesting size-‐ dependent properties, that the 5 nm β-‐FeOOH has more serious kinetic restrictions, and thus lower capacities, while it shows better cycling stability. Origins of the large overpetential have been suggested to be mainly due to the features of the surface concentration dependent potential and the slow diffusion of Na+; and the presence of the SEI layers may also contribute to the overpotential.



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ASSOCIATED CONTENT Supporting Information. HRTEM images of the prepared materials, XRD patterns of the prepared β-‐FeOOH and the materials after dis-‐ charge/charge, cycling performance of the 53 material at a -‐1 current of 500 mAh g , 1st derivative plots of the Fe K-‐edge XANES spectra in Figure 2A, EXAFS spectra of pristine β-‐ FeOOH, Fe reference and the electrode after the first dis-‐ charge, FESEM images of fresh electrode and electrode after discharge/charge, capacity data for graphite and Super P carbon electrode, and EIS spectra of electrodes at different sodiation states. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author [email protected]; [email protected]‐star.edu.sg

Funding Sources This work was supported by MOE Tier 1 Grants (RGT8/13, RG13/13, and RG131/14) of Singapore, and the Singapore Na-‐ tional Research Foundation under its Campus for Research Excellence And Technological Enterprise (CREATE) pro-‐ gramme. P.Y. is supported by SSLS via NUS Core Support C-‐ 380-‐003-‐003-‐001.

ACKNOWLEDGMENT Authors thank the Facility for Analysis, Characterisation, Testing and Simulation (FACTS) in Nanyang Technological University for materials characterizations. Authors also thank the facility support provided by the Singapore Syn-‐ chrotron Light Source.

REFERENCES (1) Slater, M. D.; Kim, D.; Lee, E.; Johnson, C. S., Sodium-‐ Ion Batteries. Adv. Funct. Mater. 2013, 23, 947-‐958. (2) Pan, H. L.; Hu, Y. S.; Chen, L. Q., Room-‐Temperature Stationary Sodium-‐Ion Batteries for Large-‐Scale Electric Energy Storage. Energy Environ. Sci. 2013, 6, 2338-‐2360. (3) Yu, H.; Ren, Y.; Xiao, D.; Guo, S.; Zhu, Y.; Qian, Y.; Gu, L.; Zhou, H., An Ultrastable Anode for Long-‐Life Room-‐ Temperature Sodium-‐Ion Batteries. Angew. Chem. Int. Ed. 2014, 53, 8963-‐8969. (4) Wang, S.; Wang, L.; Zhu, Z.; Hu, Z.; Zhao, Q.; Chen, J., All Organic Sodium-‐Ion Batteries with Na4C8H2O6. Angew. Chem. Int. Ed. 2014, 53, 5892-‐5896. (5) Billaud, J.; Clément, R. J.; Armstrong, A. R.; Canales-‐ Vázquez, J.; Rozier, P.; Grey, C. P.; Bruce, P. G., β-‐NaMnO2: A High-‐Performance Cathode for Sodium-‐Ion Batteries. J. Am. Chem. Soc. 2014, 136, 17243-‐17248. (6) Hong, S. Y.; Kim, Y.; Park, Y.; Choi, A.; Choi, N. S.; Lee, K. T., Charge Carriers in Rechargeable Batteries: Na Ions vs. Li Ions. Energy Environ. Sci. 2013, 6, 2067-‐2081. (7) Zhou, T.; Pang, W. K.; Zhang, C.; Yang, J.; Chen, Z.; Liu, H. K.; Guo, Z., Enhanced Sodium-‐Ion Battery Performance by Structural Phase Transition from Two-‐Dimensional Hexagonal-‐ SnS2 to Orthorhombic-‐SnS. ACS Nano 2014. (8) Kim, H.; Hong, J.; Park, Y.-‐U.; Kim, J.; Hwang, I.; Kang, K., Sodium Storage Behavior in Natural Graphite using Ether-‐ based Electrolyte Systems. Adv. Funct. Mater. 2015, 25, 534-‐541. (9) David, L.; Singh, G., Reduced Graphene Oxide Paper Electrode: Opposing Effect of Thermal Annealing on Li and Na Cyclability. J. Phys. Chem. C 2014, 118, 28401-‐28408.

Page 8 of 11

(10) Li, Y.; Xu, S.; Wu, X.; Yu, J.; Wang, Y.; Hu, Y.-‐S.; Li, H.; Chen, L.; Huang, X., Amorphous Monodispersed Hard Carbon Micro-‐Spherules Derived from Biomass as a High Performance Negative Electrode Material for Sodium-‐Ion Batteries. J. Mater. Chem. A 2015, 3, 71-‐77. (11) Xia, X.; Dahn, J. R., Study of the Reactivity of Na/Hard Carbon with Different Solvents and Electrolytes. J. Electrochem. Soc. 2012, 159, A515-‐A519. (12) Tang, K.; Fu, L. J.; White, R. J.; Yu, L. H.; Titirici, M.-‐ M.; Antonietti, M.; Maier, J., Hollow Carbon Nanospheres with Superior Rate Capability for Sodium-‐Based Batteries. Adv. Energy Mater. 2012, 2, 873-‐877. (13) Komaba, S.; Ishikawa, T.; Yabuuchi, N.; Murata, W.; Ito, A.; Ohsawa, Y., Fluorinated Ethylene Carbonate as Electrolyte Additive for Rechargeable Na Batteries. ACS Appl. Mater. Interfaces 2011, 3, 4165-‐4168. (14) Stevens, D. A.; Dahn, J. R., The Mechanisms of Lithium and Sodium Insertion in Carbon Materials. J. Electrochem. Soc. 2001, 148, A803-‐A811. (15) Senguttuvan, P.; Rousse, G.; Seznec, V.; Tarascon, J.-‐ M.; Palacin, M. R., Na2Ti3O7: Lowest Voltage Ever Reported Oxide Insertion Electrode for Sodium Ion Batteries. Chem. Mater. 2011, 23, 4109-‐4111. (16) Zou, W.; Li, J.; Deng, Q.; Xue, J.; Dai, X.; Zhou, A.; Li, J., Microspherical Na2Ti3O7 Prepared by Spray-‐Drying Method as Anode Material for Sodium-‐Ion Battery. Solid State Ionics 2014, 262, 192-‐196. (17) Rousse, G.; Arroyo-‐de Dompablo, M. E.; Senguttuvan, P.; Ponrouch, A.; Tarascon, J.-‐M.; Palacín, M. R., Rationalization of Intercalation Potential and Redox Mechanism for A2Ti3O7 (A = Li, Na). Chem. Mater. 2013, 25, 4946-‐4956. (18) Pan, H.; Lu, X.; Yu, X.; Hu, Y.-‐S.; Li, H.; Yang, X.-‐Q.; Chen, L., Sodium Storage and Transport Properties in Layered Na2Ti3O7 for Room-‐Temperature Sodium-‐Ion Batteries. Adv. Energy Mater. 2013, 3, 1186-‐1194. (19) Yu, X.; Pan, H.; Wan, W.; Ma, C.; Bai, J.; Meng, Q.; Ehrlich, S. N.; Hu, Y.-‐S.; Yang, X.-‐Q., A Size-‐Dependent Sodium Storage Mechanism in Li4Ti5O12 Investigated by a Novel Characterization Technique Combining in Situ X-‐ray Diffraction and Chemical Sodiation. Nano Lett. 2013, 13, 4721-‐4727. (20) Sun, Y.; Zhao, L.; Pan, H.; Lu, X.; Gu, L.; Hu, Y.-‐S.; Li, H.; Armand, M.; Ikuhara, Y.; Chen, L.; Huang, X., Direct Atomic-‐ Scale Confirmation of Three-‐Phase Storage Mechanism in Li4Ti5O12 Anodes for Room-‐Temperature Sodium-‐Ion Batteries. Nat. Commun. 2013, 4, 1870. (21) Liu, J.; Tang, K.; Song, K.; van Aken, P. A.; Yu, Y.; Maier, J., Tiny Li4Ti5O12 Nanoparticles Embedded in Carbon Nanofibers as High-‐Capacity and Long-‐Life Anode Materials for Both Li-‐Ion and Na-‐Ion Batteries. Phys. Chem. Chem. Phys. 2013, 15, 20813-‐20818. (22) Zhao, L.; Pan, H.-‐L.; Hu, Y.-‐S.; Li, H.; Chen, L.-‐Q., Spinel Lithium Titanate (Li4Ti5O12) as Novel Anode Material for Room-‐Temperature Sodium-‐Ion Battery. Chinese Phys. B 2012, 21, 028201. (23) Yu, P.; Li, C.; Guo, X., Sodium Storage and Pseudocapacitive Charge in Textured Li4Ti5O12 Thin Films. J. Phys. Chem. C 2014, 118, 10616-‐10624. (24) Wang, Y.; Yu, X.; Xu, S.; Bai, J.; Xiao, R.; Hu, Y.-‐S.; Li, H.; Yang, X.-‐Q.; Chen, L.; Huang, X., A Zero-‐Strain Layered Metal Oxide as the Negative Electrode for Long-‐Life Sodium-‐Ion Batteries. Nat. Commun. 2013, 4, 2365. (25) Wu, L.; Buchholz, D.; Bresser, D.; Gomes Chagas, L.; Passerini, S., Anatase TiO2 Nanoparticles for High Power Sodium-‐Ion Anodes. J. Power Sources 2014, 251, 379-‐385. (26) Wu, L.; Bresser, D.; Buchholz, D.; Giffin, G.; Castro, C. R.; Ochel, A.; Passerini, S., Unfolding the Mechanism of Sodium


Page 9 of 11

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Insertion in Anatase TiO2 Nanoparticles. Adv. Energy Mater. 2014, 5, 1401142. (27) Xu, Y.; Lotfabad, E. M.; Wang, H. L.; Farbod, B.; Xu, Z. W.; Kohandehghan, A.; Mitlin, D., Nanocrystalline Anatase TiO2: A New Anode Material for Rechargeable Sodium Ion Batteries. Chem. Commun. 2013, 49, 8973-‐8975. (28) Xiong, H.; Slater, M. D.; Balasubramanian, M.; Johnson, C. S.; Rajh, T., Amorphous TiO2 Nanotube Anode for Rechargeable Sodium Ion Batteries. J. Phys. Chem. Lett. 2011, 2, 2560-‐2565. (29) Li, Z.; Young, D.; Xiang, K.; Carter, W. C.; Chiang, Y. M., Towards High Power High Energy Aqueous Sodium-‐Ion Batteries: The NaTi2(PO4)3/Na0.44MnO2 System. Adv. Energy Mater. 2013, 3, 290-‐294. (30) Sun Il, P.; Gocheva, I.; Okada, S.; Yamaki, J. i., Electrochemical Properties of NaTi2(PO4)3 Anode for Rechargeable Aqueous Sodium-‐Ion Batteries. J. Electrochem. Soc. 2011, 158, A1067-‐A1070. (31) Wu, X. Y.; Cao, Y. L.; Ai, X. P.; Qian, J. F.; Yang, H. X., A Low-‐Cost and Environmentally Benign Aqueous Rechargeable Sodium-‐Ion Battery Based on NaTi2(PO4)3-‐Na2NiFe(CN)6 Intercalation Chemistry. Electrochem. Commun. 2013, 31, 145-‐148. (32) Komaba, S.; Matsuura, Y.; Ishikawa, T.; Yabuuchi, N.; Murata, W.; Kuze, S., Redox Reaction of Sn-‐Polyacrylate Electrodes in Aprotic Na Cell. Electrochem. Commun. 2012, 21, 65-‐68. (33) Lin, Y.-‐M.; Abel, P. R.; Gupta, A.; Goodenough, J. B.; Heller, A.; Mullins, C. B., Sn–Cu Nanocomposite Anodes for Rechargeable Sodium-‐Ion Batteries. ACS Appl. Mater. Interfaces 2013, 5, 8273-‐8277. (34) Liu, Y.; Xu, Y.; Zhu, Y.; Culver, J. N.; Lundgren, C. A.; Xu, K.; Wang, C., Tin-‐Coated Viral Nanoforests as Sodium-‐Ion Battery Anodes. ACS Nano 2013, 7, 3627-‐34. (35) Oh, S.-‐M.; Myung, S.-‐T.; Jang, M.-‐W.; Scrosati, B.; Hassoun, J.; Sun, Y.-‐K., An Advanced Sodium-‐Ion Rechargeable Battery Based on a Tin-‐Carbon Anode and a Layered Oxide Framework Cathode. Phys. Chem. Chem. Phys. 2013, 15, 3827-‐ 3833. (36) Darwiche, A.; Marino, C.; Sougrati, M. T.; Fraisse, B.; Stievano, L.; Monconduit, L., Better Cycling Performances of Bulk Sb in Na-‐Ion Batteries Compared to Li-‐Ion Systems: An Unexpected Electrochemical Mechanism. J. Am. Chem. Soc. 2012, 134, 20805-‐20811. (37) Qian, J.; Chen, Y.; Wu, L.; Cao, Y.; Ai, X.; Yang, H., High Capacity Na-‐Storage and Superior Cyclability of Nanocomposite Sb/C Anode for Na-‐Ion Batteries. Chem. Commun. 2012, 48, 7070-‐7072. (38) Baggetto, L.; Ganesh, P.; Sun, C.-‐N.; Meisner, R. A.; Zawodzinski, T. A.; Veith, G. M., Intrinsic Thermodynamic and Kinetic Properties of Sb Electrodes for Li-‐Ion and Na-‐Ion Batteries: Experiment and Theory. J. Mater. Chem. A 2013, 1, 7985-‐7994. (39) Zhu, Y.; Han, X.; Xu, Y.; Liu, Y.; Zheng, S.; Xu, K.; Hu, L.; Wang, C., Electrospun Sb/C Fibers for a Stable and Fast Sodium-‐Ion Battery Anode. ACS Nano 2013, 7, 6378-‐6386. (40) He, M.; Kravchyk, K.; Walter, M.; Kovalenko, M. V., Monodisperse Antimony Nanocrystals for High-‐Rate Li-‐ion and Na-‐ion Battery Anodes: Nano versus Bulk. Nano Lett. 2014, 14, 1255-‐1262. (41) Ji, L.; Gu, M.; Shao, Y.; Li, X.; Engelhard, M. H.; Arey, B. W.; Wang, W.; Nie, Z.; Xiao, J.; Wang, C.; Zhang, J.-‐G.; Liu, J., Controlling SEI Formation on SnSb-‐Porous Carbon Nanofibers for Improved Na Ion Storage. Adv. Mater. 2014, 26, 2901-‐2908. (42) Su, D.; Ahn, H.-‐J.; Wang, G., SnO2@Graphene Nanocomposites as Anode Materials for Na-‐Ion Batteries with Superior Electrochemical Performance. Chem. Commun. 2013, 49, 3131-‐3133.

(43) Wang, Y.; Su, D.; Wang, C.; Wang, G., SnO2@MWCNT Nanocomposite as a High Capacity Anode Material for Sodium-‐ Ion Batteries. Electrochem. Commun. 2013, 29, 8-‐11. (44) Wu, L.; Hu, X.; Qian, J.; Pei, F.; Wu, F.; Mao, R.; Ai, X.; Yang, H.; Cao, Y., A Sn-‐SnS-‐C Nanocomposite as Anode Host Materials for Na-‐Ion Batteries. J. Mater. Chem. A 2013, 1, 7181-‐ 7184. (45) Qu, B.; Ma, C.; Ji, G.; Xu, C.; Xu, J.; Meng, Y. S.; Wang, T.; Lee, J. Y., Layered SnS2-‐Reduced Graphene Oxide Composite – A High-‐Capacity, High-‐Rate, and Long-‐Cycle Life Sodium-‐Ion Battery Anode Material. Adv. Mater. 2014, 26, 3854-‐3859. (46) Xie, X.; Su, D.; Chen, S.; Zhang, J.; Dou, S.; Wang, G., SnS2 Nanoplatelet@Graphene Nanocomposites as High-‐Capacity Anode Materials for Sodium-‐Ion Batteries. Chem. Asian J. 2014, 9, 1611-‐1617. (47) Sun, Q.; Ren, Q.-‐Q.; Li, H.; Fu, Z.-‐W., High capacity Sb2O4 Thin Film Electrodes for Rechargeable Sodium Battery. Electrochem. Commun. 2011, 13, 1462-‐1464. (48) Yu, D. Y. W.; Prikhodchenko, P. V.; Mason, C. W.; Batabyal, S. K.; Gun, J.; Sladkevich, S.; Medvedev, A. G.; Lev, O., High-‐Capacity Antimony Sulphide Nanoparticle-‐Decorated Graphene Composite as Anode for Sodium-‐Ion Batteries. Nat. Commun. 2013, 4, 2922. (49) Kim, Y.; Park, Y.; Choi, A.; Choi, N.-‐S.; Kim, J.; Lee, J.; Ryu, J. H.; Oh, S. M.; Lee, K. T., An Amorphous Red Phosphorus/Carbon Composite as a Promising Anode Material for Sodium Ion Batteries. Adv. Mater. 2013, 25, 3045-‐3049. (50) Qian, J.; Wu, X.; Cao, Y.; Ai, X.; Yang, H., High Capacity and Rate Capability of Amorphous Phosphorus for Sodium Ion Batteries. Angew. Chem. Int. Ed. 2013, 52, 4633-‐4636. (51) Li, W.-‐J.; Chou, S.-‐L.; Wang, J.-‐Z.; Liu, H.-‐K.; Dou, S.-‐ X., Simply Mixed Commercial Red Phosphorus and Carbon Nanotube Composite with Exceptionally Reversible Sodium-‐Ion Storage. Nano Lett. 2013, 13, 5480-‐5484. (52) Li, W.; Chou, S.-‐L.; Wang, J.-‐Z.; Kim, J. H.; Liu, H.-‐K.; Dou, S.-‐X., Sn4+xP3 @ Amorphous Sn-‐P Composites as Anodes for Sodium-‐Ion Batteries with Low Cost, High Capacity, Long Life, and Superior Rate Capability. Adv. Mater. 2014, 26, 4037-‐4042. (53) Qian, J.; Xiong, Y.; Cao, Y.; Ai, X.; Yang, H., Synergistic Na-‐Storage Reactions in Sn4P3 as a High-‐Capacity, Cycle-‐stable Anode of Na-‐Ion Batteries. Nano Lett. 2014, 14, 1865-‐1869. (54) Kim, Y.; Kim, Y.; Choi, A.; Woo, S.; Mok, D.; Choi, N.-‐ S.; Jung, Y. S.; Ryu, J. H.; Oh, S. M.; Lee, K. T., Tin Phosphide as a Promising Anode Material for Na-‐Ion Batteries. Adv. Mater. 2014, 26, 4139-‐4144. (55) Fullenwarth, J.; Darwiche, A.; Soares, A.; Donnadieu, B.; Monconduit, L., NiP3: A Promising Negative Electrode for Li-‐ and Na-‐Ion Batteries. J. Mater. Chem. A 2014, 2, 2050-‐2059. (56) Kovnir, K. A.; Kolen’ko, Y. V.; Ray, S.; Li, J.; Watanabe, T.; Itoh, M.; Yoshimura, M.; Shevelkov, A. V., A Facile High-‐Yield Solvothermal Route to Tin Phosphide Sn4P3. J. Solid State Chem. 2006, 179, 3756-‐3762. (57) Fox, M., Dangerous When Wet Materials and Division 4.3. In Glossary for the Worldwide Transportation of Dangerous Goods and Hazardous Materials, Springer Berlin Heidelberg: 1999; pp 58-‐60. (58) Longwell, J. P., Alternative Technologies for the Destruction of Chemical Agents and Munitions. National Academy Press: 1993. (59) Yu, L. H.; Xi, S. B.; Wei, C.; Zhang, W. Y.; Du, Y. H.; Yan, Q. Y.; Xu, Z. C., Superior Lithium Storage Properties of β-‐ FeOOH. Adv. Energy Mater. 2015, 5, 1401517. (60) Cabana, J.; Monconduit, L.; Larcher, D.; Rosa Palacin, M., Beyond Intercalation-‐Based Li-‐Ion Batteries: The State of the Art and Challenges of Electrode Materials Reacting Through Conversion Reactions. Adv. Mater. 2010, 22, E170-‐E192.



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(61) Huang, B.; Tai, K.; Zhang, M.; Xiao, Y.; Dillon, S. J., Comparative Study of Li and Na Electrochemical Reactions with Iron Oxide Nanowires. Electrochim. Acta 2014, 118, 143-‐149. (62) Oh, S.-‐M.; Myung, S.-‐T.; Yoon, C. S.; Lu, J.; Hassoun, J.; Scrosati, B.; Amine, K.; Sun, Y.-‐K., Advanced Na[Ni0.25Fe0.5Mn0.25]O2/C–Fe3O4 Sodium-‐Ion Batteries Using EMS Electrolyte for Energy Storage. Nano Lett. 2014, 14, 1620-‐1626. (63) Valvo, M.; Lindgren, F.; Lafont, U.; Bjorefors, F.; Edstrom, K., Towards More Sustainable Negative Electrodes in Na-‐Ion Batteries via Nanostructured Iron Oxide. J. Power Sources 2014, 245, 967-‐978. (64) Hariharan, S.; Saravanan, K.; Ramar, V.; Balaya, P., A Rationally Designed Dual Role Anode Material for Lithium-‐Ion and Sodium-‐Ion Batteries: Case Study of Eco-‐Friendly Fe3O4. Phys. Chem. Chem. Phys. 2013, 15, 2945-‐2953. (65) Jian, Z.; Zhao, B.; Liu, P.; Li, F.; Zheng, M.; Chen, M.; Shi, Y.; Zhou, H., Fe2O3 Nanocrystals Anchored onto Graphene Nanosheets as the Anode Material for Low-‐Cost Sodium-‐Ion Batteries. Chem. Commun. 2014, 50, 1215-‐1217. (66) Zhang, N.; Han, X.; Liu, Y.; Hu, X.; Zhao, Q.; Chen, J., 3D Porous γ-‐Fe2O3@C Nanocomposite as High-‐Performance Anode Material of Na-‐Ion Batteries. Adv. Energy Mater. 2014, 1401123.

Page 10 of 11

(67) Lopez, M. C.; Lavela, P.; Ortiz, G. F.; Tirado, J. L., Transition Metal Oxide Thin Films With Improved Reversibility as Negative Electrodes for Sodium-‐Ion Batteries. Electrochem. Commun. 2013, 27, 152-‐155. (68) Jiang, Y.; Hu, M.; Zhang, D.; Yuan, T.; Sun, W.; Xu, B.; Yan, M., Transition Metal Oxides for High Performance Sodium Ion Battery Anodes. Nano Energy 2014, 5, 60-‐66. (69) Taberna, L.; Mitra, S.; Poizot, P.; Simon, P.; Tarascon, J. M., High Rate Capabilities Fe3O4-‐Based Cu Nano-‐Architectured Electrodes for Lithium-‐Ion Battery Applications. Nat. Mater. 2006, 5, 567-‐573. (70) Yu, L. H.; Wei, C.; Yan, Q. Y.; Xu, Z. C., Controlled Synthesis of High-‐Performance β-‐FeOOH Anodes for Lithium-‐ Ion Batteries and Their Size Effects. Nano Energy 2015, 13, 397-‐ 404. (71) Delmer, O.; Balaya, P.; Kienle, L.; Maier, J., Enhanced Potential of Amorphous Electrode Materials: Case Study of RuO2. Adv. Mater. 2008, 20, 501-‐505. (72) Conway, B. E., Transition from “Supercapacitor” to “Battery” Behavior in Electrochemical Energy Storage. J. Electrochem. Soc. 1991, 138, 1539-‐1548.


Page 11 of 11

3.0

1st 2nd 3rd 4th 5th

2.5 2.0 1.5 1.0 0.5 0.0 0

200 400 600 800 Capacity (mAh g-1)


Authors are required to submit a graphic entry for the Table of Contents (TOC) that, in conjunction with the man-‐ uscript title, should give the reader a representative idea of one of the following: A key structure, reaction, equation, concept, or theorem, etc., that is discussed in the manuscript. Consult the journal’s Instructions for Authors for TOC graphic specifications.

Potential (V, vs. Na +/Na)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


1st discharge

(0) Fe (III) 7120

β-FeOOH 1.1 V

0.8 V 0.3 V 0.005 V Fe

7140 7160 Energy (eV)

Insert Table of Contents artwork here


11

Î²-FeOOH: An Earth-Abundant High-Capacity Negative Electrode

Recommend Documents